SPRAY ROLLING OF RAPIDLY SOLIDIFIED Al-Fe-V-Si ALLOY

A THESIS SUBMITTED TO THE GRADUATE SCHOOL OF NATURAL AND APPLIED SCIENCES

OF MIDDLE EAST TECHNICAL UNIVERSITY

BY

HALİT AKIN ÖZYURDA

IN PARTIAL FULFILLMENT OF THE REQUIREMENTS FOR

THE DEGREE OF MASTER OF SCIENCE IN

METALLURGICAL AND MATERIALS ENGINEERING

MAY 2006

Approval of the Graduate School of Natural and Applied Sciences

Prof. Dr. Canan ÖZGEN

Director

I certify that this thesis satisfies all the requirements as a thesis for the degree of Master of Science

Prof. Dr. Tayfur ÖZTÜRK

Head of Department

This is to certify that we have read this thesis and that in our opinion it is fully adequate, in scope and quality, as a thesis for the degree of Master of Science. Prof. Dr. Ali KALKANLI Supervisor Examining Committee Members Prof. Dr. Rıza GÜRBÜZ (METU, METE)

Prof. Dr. Ali KALKANLI (METU, METE)

Prof. Dr. Cevdet KAYNAK (METU, METE)

Assoc. Prof. Dr. C. Hakan GÜR (METU, METE)

Prof. Dr. Tamer ÖZDEMİR (GAZI UNV.)

iii

I hereby declare that all information in this document has been obtained and presented in accordance with academic rules and ethical conduct. I also declare that, as required by these rules and conduct, I have fully cited and referenced all material and results that are not original to this work.

Name, Last name : Halit Akın ÖZYURDA

Signature :

iv

ABSTRACT

SPRAY ROLLING OF RAPIDLY SOLIDIFIED Al-Fe-V-Si ALLOY

ÖZYURDA, Halit Akın

Ms.Sc. Department of Metallurgical and Materials Engineering

Supervisor: Prof. Dr. Ali KALKANLI

May 2006, 78 pages In this study an experimental spray-rolling set-up is designed in order to produce

rapidly solidified Al-Fe-V-Si flat product. Al-Fe-V-Si alloys produced by rapid

solidification powder metallurgy (RSP/M) methods are mostly used in high

temperature applications in aerospace and automotive industries. The RSP/M

technique used is spray deposition, which is desirable because of the high cooling

rates achieved, as a result fine silicide dispersoids and intermetallics are observed

in the microstructure which are known to contribute to the mechanical properties

i.e. high strength at elevated temperatures, thermal stability, fracture toughness,

corrosion resistance. Since spray deposition is a droplet consolidation process a

considerable amount of porosity is expected in the final product. In this work,

spray rolling process, which consists of spray deposition and subsequent hot twin-

rolling stage, is designed and developed by interpreting the results obtained from

SEM, XRD, tensile, three point bending and hardness tests of the specimens

formed in several design stages. Two original intermetallic phases characterized in

this study are V3Si and V2Mg3Al18 .

Keywords: AlFeVSi alloy, spray rolling method, rapid solidification, spray deposition

v

ÖZ

HIZLI KATILAŞMIŞ Al-Fe-V-Si ALAŞIMININ PÜSKÜRTME HADDELENMESİ

ÖZYURDA, Halit Akın

Yüksek Lisans, Metalurji ve Malzeme Mühendisliği Bölümü

Tez Yöneticisi: Prof. Dr. Ali KALKANLI

Mayıs 2006, 78 sayfa

Bu çalışmada hızlı katılaşmış düz şerit halinde AlFeVSi alaşımı ürününün

üretilmesini hedefleyen deneysel bir püskürtme haddeleme düzeneği

tasarlanmaktadır. Hızlı katılaşma toz metalürjisi (HKT/M) yöntemleriyle üretilen

AlFeVSi alaşımları ağırlıkla uçak ve otomotiv endüstrisinde yüksek sıcaklıktaki

uygulamalarda kullanılırlar. Bu çalışmada HKT/M yöntemi olarak püskürtme

biriktirme kullanılmıştır. Püskürtme biriktirme yöntemi, yüksek sıcaklıklardaki

güç, termal kararlılık, kırılma tokluğu, korozyon dayancı gibi mekanik özelliklere

katkıda bulunduğu bilinen ince silisid tanecikleri ve intermetaliklerin içyapıda

gözlenmesini sağlayan yüksek soğuma hızlarına ulaşılabildiği için istenmiştir.

Püskürtme biriktirmenin bir damlacık yoğunlaştırma yöntemi olmasından dolayı

son üründe gözeneklilik beklenmektedir. Bu çalışmada, püskürtme biriktirme

sürecini takiben bir sıcak çift haddeleme aşamasından oluşan püskürtme

haddeleme süreci, tasarım aşamalarından elde edilen numunelerin SEM, XRD,

çekme, üç nokta bükme ve sertlik değerlendirmelerinden elde edilen sonuçların

yorumlanması ile tasarlanacak ve geliştirilecektir. Bu çalışmada tanımlanan özgün

intermetallic fazları V3Si ve V2Mg3Al18 ‘dir.

Anahtar Kelimeler: AlFeVSi alaşımı, püskürtme haddeleme yöntemi, hızlı katılaşma, püskürtme döküm

vi

To My Family

vii

ACKNOWLEDGEMENTS

I would like to express my deepest gratitude and appreciation to my supervisor

Prof. Dr. Ali Kalkanlı for his support, inspiration, motivation, criticism and

excellent guidance since my graduation project.

I would also like to express my great thanks my co-worker Kağan Menekşe for his

support, ideas and friendship.

I would like to express my special thanks to the Mastırkastır Team (Kağan

Menekşe, Arda Çetin, Salih Türe and Ufuk Cengizel) for their assistance and

friendship through the experimental work of this study.

I would like express my sincere thanks to my parents; Ümit - Ferda Özyurda and

my sister Deniz for their never-ending love, support, motivation and guidance

during this study and during my life.

Finally, I would like to thank Sahir Gül, Mehmet Oğuz and Müfit Özyurda for

their support, understanding and motivation during this study.

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TABLE OF CONTENTS

PLAGIARISM.........................................................................................................iii

ABSTRACT.............................................................................................................iv

ÖZ.............................................................................................................................v

DEDICATION.........................................................................................................vi

ACKNOWLEDGEMENTS………………………………………………………vii

TABLE OF CONTENTS.......................................................................................viii

CHAPTERS

1. INTRODUCTION.................................................................................………..1

2. THEORY AND LITERATURE SURVEY........................................................3

2.1 Elevated Temperature Aluminum Alloys.................................................... 3

2.2 Rapidly Solidified Al-Fe-V-Si Alloy.......................................................…. 5

2.3 Spray Deposition.......................................................................................... 7

2.3.1 Features and History of Spray Deposition........................................... 7

2.3.2 Fundamentals of Spray Deposition.....................................................10

2.4 Spray Rolling............…...............................................................................14

3. EXPERIMENTAL PROCEDURE................................................................... 17

3.1 Alloy Preparation and Composition.......................…................................. 17

3.2 Spray Rolling Set-up Design........................................................................ 18

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3.2.1 Early Spray Rolling Setups’ Design Stages....……..........................20

3.2.2 Advanced Design of Spray Rolling Setup…………………………25

3.3 Hot Rolling………………………………………………………………32

3.4 X-Ray Diffraction Study………………………………………………...32

3.5 Optical and Scanning Electron Microscopy Study……………………...32

3.6 Hardness Test……………………………………………………………33

3.7 Three Point Bending Test……………………………………….…….... 33

3.8 Tensile Test……………………………………………….………...……34

3.9 True Density Measurement Study…………………………………...…..34

4. RESULTS AND DISCUSSION..................................................................... 35

4.1 Phase Characterization...............................................................................35

4.1.1 Optical and Scanning Electron Microscopy Results………………..36

4.1.2 X-Ray Diffraction Results…………………………………………. 54

4.2 Mechanical Test Results………………………………………………… 58

4.2.1 Three Point Bending Test Results…………………………………. 58

4.2.2 Tensile Test Results…………………………………………………60

4.2.3 Hardness Test Results……………………………………………….61

4.3 True Density Measurement Results……………………………………...63

5. CONCLUSIONS......................................................................................….. 64

6. SUGGESTIONS FOR FUTURE WORK………………………………….. 66

REFERENCES................................................................................................... 67

APPENDICES.................................................................................................... 72

A: X-RAY DIFFRACTION CARDS OF THE CHARACTERIZED PHASES. 72 B: SAMPLE STRESS VS STRAIN GRAPH OF B2 SPECIMEN.................... .78

1

CHAPTER1

INTRODUCTION

In this study a 8009 series aluminum alloy is concerned which has the alloying

elements Fe, V and Si. This alloy is designed by Allied Signal Inc. in the 1980s

[6]. The alloy is designed by the addition of V to the Al-Fe-Si ternary alloy, the

vanadium addition showed that when processed by rapid solidification a fine

silicide dispersiod phase is stabilized in the microstructure giving the alloy

exceptional room and elevated temperature properties.

The rapid solidification processing which is necessary in production of Al-Fe-V-Si

alloys is mostly applied by powder metallurgy routes including the powder

production and consolidation steps. The main objective of this study is to assess

the production of the Al-Fe-V-Si with a different process than the conventional

powder metallurgy route, namely the spray rolling which consists of two main

stages spray deposition of the molten alloy for rapid solidification and subsequent

hot rolling for the simultaneous consolidation of the end product in the same

process. This process is firstly introduced by A.R.E. Singer in 1970s [18,19,22]

and substantially developed until present. The attempt to use the spray rolling

process for the production of Al-Fe-V-Si alloy is believed to result in valuable data

and contribution when the properties of related alloy resulting from rapid

solidification is concerned.

2

The process design of spray rolling of Al-Fe-V-Si alloy is carried out by control of

the process parameters together with the interpreting of the results obtained from

the process design stages by the usage of testing and characterization tools as

scanning electron microscopy, X-ray diffraction and mechanical tests like tensile,

three point bending, hardness. The conclusions derived from the analysis of the

process by assessment of results obtained from the mentioned tools, are used for

evolution of the laboratory scale experimental spray-rolling set-up.

3

CHAPTER 2

THEORY AND LITERATURE SURVEY

2.1 Elevated Temperature Aluminum Alloys

Powder metallurgy technology is inherently appropriate for designing alloys for

elevated temperature service. Rapid solidification conditions attained by powder

metallurgy technology allows the formation of finely dispersed intermetallic

strengthening phases which give resistance to coarsening and are not practical, or

in some situations not possible to be formed by conventional ingot metallurgy

methods.

Early studies on the elevated temperature aluminum alloy design were aiming the

service temperature as 315 to 345°C. But in 1980s U.S. Air Force extended the

possible service temperature to a challenging state of 480°C. Much of the

development in this area is achieved by dispersing slow-diffusivity transition

metals as intermetallic phases in aluminum and most of the alloys developed

contain iron. [1]

Scientists of Alcoa investigated numerous alloys and they designed the

commercial Al-Fe-Ce alloys for elevated temperature service with the designations

of CU78 and CZ42. These alloys show good tensile strength up to 315°C and also

good room temperature properties after exposed to elevated temperatures. [2,3,4] It

is also reported that these alloys also has good resistance to environmentally

assisted cracking. [4]

4

The leading RS-P/M elevated temperature aluminum alloys developed are the Al-

Fe-V-Si alloys designed by Allied Signal Inc., which are produced by planar flow

casting [5,6]. Alloys are named due to their alloying elements as FVS0812,

FVS1212 and FVS0611. FVS0812 and FVS1212 showed very high strengths at

315°C and usable strengths at 425°C. It is reported that the microstructures of

these alloys have very finely dispersed silicides that contribute greatly to the

strength over a wide temperature range without degrading the corrosion resistance.

FVS0812 alloy is used for aircraft wheels and replaced the 2014 alloy. Moreover,

its elevated temperature stability and high strength at temperatures above 300°C

makes it challenging to titanium alloys used in some engine components.

Allied Signal Inc. produces the FVS0611 alloy for applications that require good

formability, which has a lower alloying element content.

Another elevated temperature aluminum alloy is developed by scientists in Alcan

Int. Ltd. They used the Al-Cr-Zr system because of the lower sensitivity to the

cooling rate and the better hot workability. It is stated that both chromium and

zirconium can produce thermally stable solid solutions even at modest

solidification rates. After the consolidation of the powder product, it could be

extruded or rolled with low stresses and they aging could be done to obtain stable

intermetallics, which give the elevated temperature strength. [7]

5

Table 2.1 Mechanical properties of elevated temperature aluminum alloys. [1]

Material Temperature

°C °F Yield Strength MPa ksi

Tensile Strength MPa ksi

Elongation %

Fracture Toughness

KIC

MPa√m ksi√in.

Al-8Fe-7Ce 25 77 316 600

418.9 60.8 178.1 25.8

484.9 70.3 193.8 28.1

7.0 7.6

8.5 7.7 7.9 7.2

Al-8Fe-2Mo-1V 25 77 316 600

323.5 46.9 170.0 24.6

406.6 59.0 187.5 27.2

6.7 7.2

9.0 8.2 8.1 7.4

Al-10.5Fe-2.5V 25 77 316 600

464.1 67.3 206.3 29.9

524.5 76.1 240.0 34.8

4.0 6.9

5.7 5.2 8.1 7.4

Al-8Fe-1.4V-1.7Si 25 77 316 600

362.5 52.6 184.4 26.7

418.8 60.7 193.8 28.1

6.0 8.0

36.4 33.1 14.9 13.6

2.2 Rapidly Solidified Al–Fe–V–Si Alloy

In recent decades there has been a considerable amount of effort devoted to the

development of aluminum alloys with high temperature stability which can

compete with titanium alloys on a specific strength basis. Al-Fe-V-Si alloys have

attracted considerable interest and are widely used in aerospace and automotive

industry applications due to excellent room and high temperature strength, thermal

stability, fracture toughness and good corrosion resistance [8-10]. These alloys are

generally processed through rapid solidification powder metallurgy processing

routes due to low diffusivity Fe and V in Al in high cooling rates [12] and it is

observed that addition of V to the Al-Fe-Si alloy stabilizes the cubic Al12(Fe,V)3Si

phase over the detrimental Al-Fe intermetallics and also the silicide phase is said

to be responsible for the low coarsening rate up to about 773K giving the alloy its

high temperature properties[13]. It is also reported that the microstructure Al-Fe-

V-Si alloys consist of very fine, nearly spherical Al12(Fe,V)3Si (silicide)

dispersoids, formed from the decomposition of the rapidly solidified

microstructures, uniformly distributed throughout the aluminum matrix [14,15].

Moreover, these alloys are devoid of any coarse needle or plate-like intermetallic

phases that degrade alloy ductility and fracture toughness [14]. Recently, it is

6

reported that under normal die casting conditions ten-armed star shaped crystals of

Al3Fe intermetallic phase precipitates [16] and they may act as stress raisers being

detrimental to mechanical properties [17]. The Al-Fe-V-Si alloys contain high

transition element concentrations (5.5-12.5 wt% Fe) that form high volume

fractions of silicides and their desirable microstructural morphology, which can be

achievable by rapid solidification techniques, essentially controls the mechanical

properties of these alloys.

Table 2.2 Mechanical properties of Al-Fe-V-Si alloy [1].

Test Temperature °C °F

Yield Strength MPa ksi

Ultimate Tensile Strength

MPa ksi

Elongation %

Young’s Modulus Gpa 106 psi

24 75 413 60 462 67 12.9 88.4 12.8 149 300 345 50 379 55 7.2 83.2 12.0 232 450 310 45 338 49 8.2 73.1 10.6 316 600 255 37 276 40 11.9 65.5 9.5 427 800 138 20 155 22 15.1 61.4 8.9

Figure 2.1 TEM micrograph showing the fine dispersed Al13(Fe,V)3Si

silicides. [11]

7

2.3 Spray Deposition

2.3.1. Features and History of Spray Deposition

Although numerous property improvements have been demonstrated to result from

rapid solidification, commercialization has been limited due to difficulties

associated with the production of bulk shapes. These problems are largely

associated with the oxides present on aluminum particulates. In an effort to control

the volume fraction of oxides present in consolidated reactive rapidly solidified

alloys obtained by powder metallurgy (PM), and to simplify the overall

processing, spray atomization and collection processes are being studied in several

institutions.

The principles of spray deposition were pioneered during the 1970s by A. Singer

at the University College of Swansea, United Kingdom. Singer proposed the

production of rolled strip directly from molten metal as an alternative to the

current practice of casting and rolling large ingots [18,19] Singer studied the spray

rolling of metals as an alternative to a process originally developed at the Reynolds

Metal Company in 1967 [20]. In this process aluminum was centrifugally

atomized, reheated, fed into roll gaps and hot-rolled to produce strip in a

continuous operation [20-22]. The general principle of spray deposition is to

atomize a stream of molten metal by means of high velocity gases (i.e., argon or

nitrogen) and to direct the resulting spray into or onto a shaped collector (mold or

substrate). On impact with the collector the particles flatten and weld together to

form a hot, high density preform which can be readily forged to form a fully dense

product. The surface of the substrate must be prepared in such a way that the first

layer of atomized droplets will adhere to form a smooth quenched layer. Singer

conducted most of his work on aluminum alloys. The atomized droplets were

approximately 100-150µm in diameter, and were produced by using conventional

atomization equipment. The process parameters used in these studies were

sufficient to atomize a melt stream and to cool the particles to temperatures near

8

the solidus during a flight path of about 18 in (0.45 m). One of the main problems

associated with this process of strip making was that the thickness of the as-

deposited strip was not uniform across the width [23].

Several years later, Osprey Metals Ltd. (Neath, South Wales) successfully applied

the early ideas of Singer for the production of forging preforms. This effort was

undertaken as a result of increasing pressures on the metal-forming industry to

reduce manufacturing costs, improve material utilization, improve material

structure and properties, and increase efficiency [21-28]. The Osprey process

became an attractive alternative to the two existing basic methods of producing

forgings: (1) the forging route, and (2) the PM method. The conventional forging

route employing rolled stock was very inefficient on account of the large quantities

of scrap generated. The powder process is relatively expensive as a result of the

necessary sieving, reducing, blending, compacting and heating, and due to

problems associated with oxide contamination [25].

It would appear that a spray deposition approach will inherently avoid the extreme

thermal excursions, with concomitant degradation in mechanical properties and

extensive macrosegregation, normally associated with casting processes.

Furthermore, this approach also eliminates the need to handle fine reactive

particulates, as is necessary with powder metallurgical processes. A series of

studies based on the early spray atomization and collection processes, but using

high pressure gas atomization for the spraying of fine, rapidly quenched droplets

under a tightly controlled atmosphere resulted in the development of other spray

deposition processes; these include liquid dynamic compaction (LDC) [29], and

variable co-deposition of multiphase materials (VCM) [30].

One particular spray deposition set-up is shown in Figure 2.2.

9

Figure 2.2 Spray deposition set-up [31]

10

2.3.2 Fundamentals of Spray Deposition

In spray atomization and deposition process, the microstructural characteristics of

the spray-deposited material depend to a great extent on the conditions of the

droplets prior to impact; that is, on the relative proportions of liquid and solid

present, temperature, velocity, and size of the microstructure in both the partially

and fully solidified droplets. It is also important to know the size and distribution

of the droplets since the amount of heat dissipated by droplets during flight is

strongly dependent on their size. The microstructural evolution during spray

atomization and deposition can be separated into two distinct but closely related

stages. The first stage encompasses those phenomena that are primarily active in

the atomized spray prior to impact. The second stage commences after the droplets

have impacted the deposition surface, and alteration of the microstructure resulting

from impact must be considered.

At the moment of impact, the thermal and solidification conditions of the droplet

distribution will depend on;

(1) the thermodynamic properties of the material,

(2) the thermodynamic properties of the gas,

(3) the processing parameters.

The first step towards developing a better understanding of the evolution of the

microstructure after deposition and impact, is to establish the thermal conditions of

the growing deposit. Recently, various investigators computed the temperature

distribution in a growing deposit utilizing a differential thermal energy balance

[32]. In addition to unidirectional heat flow, their analysis incorporates the

following assumptions; (1) the thermal conditions of the spray during impact can

11

be approximated utilizing enthalpy considerations, (2) at the preform-substrate

interface, conduction of heat is governed by an interface heat transfer coefficient,

and (3) during deposition, the top region of the deposit loses heat through the

combined effects of convection and radiation.

The thermal history of the material throughout deposition is dictated, in part, by

the rate of transfer of thermal energy from the atomized spray into the deposited

material. In turn, the rate of transfer of thermal energy from the spray into the

deposit is directly related to the average enthalpy content of the spray at the

moment of impact. The average enthalpy content of the spray can be estimated

from the weighted average of the total heat content of each and every droplet in the

spray. Hence, the total enthalpy content can be calculated from [33].

Hspray= ΣHp(di)di3f(di)/ Σdi

3f(di)

where Hp(di) is the enthalpy content of a droplet of diameter di and f(di) is the

fraction of droplets with diameter di.

An important microstructural characteristic frequently associated with spray

deposited structures is the presence of a finite amount of non-interconnected

porosity [30,33,34]. The overall amount of porosity present in spray-deposited

materials depends on: (1) the thermodynamic properties of the material, (2) the

thermodynamic properties of the gas, and (3) the processing parameters. Under

conditions typical for aluminum alloys, the amount of porosity present in spray-

deposited materials has been reported to be in the 1-10% range. In addition, the

distribution of pore sizes has been shown to be skewed, with a large proportion of

pores in the 1-10 µm size range [33,37]. It is interesting to note that several

investigators have reported that a large proportion of the fine pores are

preferentially located at grain boundaries [30,35,36].

The origin of porosity in spray-deposited materials can be traced to three possible

12

sources: (1) gas entrapment, (2) solidification shrinkage, and (3) interstitial

porosity. The first source, entrapment during impact and solidification, is to be

expected as a result of the considerable quantities of gas used to disintegrate the

metal during atomization. Recent results, however, suggest that the development

of porosity from gas entrapment is not as important a mechanism as originally

thought. Ogata et al. [37] reported that the amount of gas contained in a spray-

deposited nickel-base superalloy atomized with argon amounted to less than 3 ppm

argon. Similarly, fast neutron activation analysis of various spray-atomized

aluminum alloys indicated nitrogen levels of approximately 5 ppm [33]. Finally, it

is shown that (1) the irregular morphology of the pores, and (2) the limited solu-

bility of the inert gases commonly used for atomization, it is highly improbable

that a large proportion of the pores originate from the expansion of entrapped

gases [35].

Owing to the limited amount of liquid phase present under proper deposition

conditions, it is highly unlikely that solidification shrinkage plays an important

role in the formation of the observed pore distribution. Hence, it is anticipated that

only under deposition conditions where there is an excessive amount of liquid

phase at impact, will this mechanism play a significant role in the formation of

porosity. It is important to note, however, that in the presence of excess amounts of

liquid phase as a result of: (1) coarse droplet sizes, and/or (2) remelting of solid

phases caused by high spray enthalpies, turbulent interactions between the

atomizing gas and the molten material will result in the formation of large amounts

of porosity. The available experimental evidence suggests that a large proportion

of the porosity present develops from interstices formed as the droplets impact on

one another, leaving micron-sized, irregular cavities as they overlap. This

mechanism is consistent with the observed correlation between deposition

conditions such as spray density, powder size, and fraction solidified, and the

amount of porosity present throughout the deposit.

An interesting attribute of spray-deposited materials is their microstructural

13

characteristics. Some investigators [30], studied the microstructure of spray-

atomized and deposited Fe-Nd-X materials and reported that although most of

structure consisted of equiaxed grains, three types of microstructures could be

distinguished: equiaxed grains, 2-20 µm; solid powders 1-5 µm; and splats 0.1-1.0

µm in thickness. The presence of these three types of microstructural features was

related to the splatting conditions during deposition. At the beginning of the

experiment, when spray density as well as the droplet size are largest, it is possible

that complete solidification of splats may not occur before the arrival of the next

partially solidified or undercooled droplets. Accordingly, a thin film of liquid will

be retained on the deposition surface. The mechanical action of the splatting will

break off any oxide layer on the droplet surface as well as the delicate dendrite

arms in both the depositing droplets and in the liquid film. As a result an equiaxed

grain morphology in the microstructure is formed.

To sum up, it can be said that spray deposited materials show characteristic

microstructural features as fine equiaxed grained microstructures, absence of

metastable phase formation, excess solid solubility, and the absence of

macrosegregation.

Finally, to effectively gain advantage of the features of spray deposition explained

above, It is reported that, quality and effectiveness of the spray deposition is

significantly affected by process parameters so in order to achieve the required

deposit shape an desired properties, the following process parameters must be

controlled and optimized for the production [38,39,40];

1. The diameter of the melt orifice, hence the melt flow rate,

2. The gas pressure of the atomizer nozzle, hence the gas flow rate,

3. The distance of the substrate from the atomizer nozzle,

4. The inclination of the spray cone due to the substrate,

5. The gas impinging angle,

6. The protrusion lenght of the nozzle,

14

7. The shape of the substrate itself,

8. The movement of the substrate.

9. The amount of superheating of the melt.

2.4 Spray Rolling

High strength aluminum alloys are used extensively for aerospace applications.

Flat products are manufactured by conventional ingot metallurgical (I/M)

processing. Ingots are direct-chill (DC) cast to about 0.6m thickness, scalped,

homogenized and hot rolled to the desired thickness. Following this, the material is

further processed (e.g. heat treated, cold rolled to final gauge, etc.) according to

temper requirements and desired properties. I/M processing remains the most

reliable, versatile production method for these alloys. However, it is energy and

capital equipment intensive, reflecting the need to homogenize ingots and hot

work casting flaws.

Twin-roll casting, was originally proposed by Bessemer in the mid 1800s [41]. It

combines solidification and hot rolling in a single operation. Liquid metal is fed

into the gap between large water-cooled hollow rolls, where it solidifies to form

strip up to about 6mm thick. Since its use began in industry, about 50 years ago,

the technology has improved steadily, particularly in the last 20 years. However,

twin-roll casters operate much slower than their theoretical production-rate limit to

satisfy quality requirements. And, due to production rate and quality issues,

commercial sheet has been limited to alloys that have a suitably narrow freezing

range [42-47].

At the present time, twin-roll casting is not used commercially to process the 2124

and 7050 alloys. A new strip/sheet casting process, termed “spray rolling” (also

“spray strip casting”), is currently under development.[48]. The general concept of

spray rolling, i.e., deposition into a roll gap to form strip, is credited to A. Singer

15

who conducted pioneering work during the 1970s [18]. In general terms, spray

rolling combines features of twin-roll casting and spray forming (spray

deposition). Spray rolling consists of atomizing molten metal with a high velocity

inert gas, extracting most of the metal’s latent heat in-flight via convective cooling

by entrained inert gas (to about 70% solid), and depositing the atomized droplets

between mill rolls. The metal is consolidated into strip/sheet while still in a semi-

solid and highly formable condition. As with twin-roll casting, it is believed that

some solid state compaction (hot rolling) occurs as the strip advances through the

roll gap. While spray rolling shares many similarities with twin-roll strip casting,

there are important differences:

1. In twin-roll casting, the metal’s superheat and latent heat are dissipated almost

exclusively by conductive heat transfer to water-cooled rolls. In spray rolling,

convective heat transfer from atomized droplets plays a prominent role and teams

with conductive transfer at the rolls to increase production rate.

2. The metal introduced to the rolls in twin-roll casting is molten, while in spray

rolling, it has a “slushy” character. Solid particles in the slush act as nucleation

sites, producing an equiaxed grain structure and limiting segregation

Aluminum alloys with high solute content and broad freezing ranges have been

successfully spray rolled. While still in the early stages of development, spray

rolling shows promise for reducing strip/sheet manufacturing costs while

improving quality. The inherent rapid solidification and solid solubility extension

may, in the future, provide an interesting avenue for the development of alloys

tailored for the process and which show unique combinations of properties.

Advantages of spray rolling over conventional I/M processing appear to include

cost reduction and the elimination of energy intensive unit operations such as ingot

casting, homogenization and hot rolling. When compared to twin-roll strip casting,

advantages appear to be the high quality and production rate, and the ability to

process a broader range of alloys[48].

16

The spray rolling set-up designed by K.M. McHugh et al. [48] is shown in the

Figure 2.3.

Figure 2.3 Spray rolling set-up designed by McHugh et al.

17

CHAPTER 3

EXPERIMENTAL PROCEDURE

3.1 Alloy Preparation and Composition

The alloy is melted in inductotherm induction furnace at 800oC and then poured

into the furnace of spray rolling set-up, which allows the flow of the alloy melt

thorough to the atomizer nozzle.

Two different compositions of Al-Fe-V-Si alloy are used in this study, alloys are

prepared from pure Al, pure Fe, ferrovanadium FeV, pure Mg and pure Si. Mg is

added to the commercial Al-Fe-V-Si alloy because of its known improvements;

including the impeding of the growth of the Al-Fe intermetallics in low cooling

rates previously studied by Sahoo et al. [49], decreasing the viscosity of the melt

experienced in previous studies and ability to form Mg2Si resulting in precipitation

hardening when heat treatment is conducted.

The alloy compositions are given in Table 3.1

18

Table 3.1 Table of alloy compositions

3.2 Spray Rolling Set-up Design

As it was stated before the main aim in this study is achieve the continuos

production of the Al-Fe-V-Si alloy strip, which is formed by spray deposition and

subsequent rolling. To achieve this, horizontal spray deposition is chosen as the

rapid solidification route and a horizantal twin-roller is integrated to the set-up.

In the spray rolling process studied in this study, air atomized molten alloy is

directed to the rolls aiming the formation of a semi solid deposit between the rolls

and subsequent solidification with simultaneous deformation during the passage

through the twin roller. The most critical stage in spray rolling process is presumed

to be the spray deposition stage since the control of the spray cone is essential.

The important process parameters that guide the design stages is summarized and

process variables are classified below.

Al(%wt) Fe (%wt) V(%wt) Si(%wt) Mg(%wt)

Alloy 1 Bal. 8 1.7 8 1.5

Alloy 2 Bal. 8 1.4 1.8 1.3

Alloy 2 has a close composition to the commercial alloy FVS812 without the Mg addition.

19

- The diameter of the melt orifice, hence the melt flow rate;

- The gas pressure of the atomizer nozzle, hence the gas flow rate

- The amount of superheating of the melt

- The distance of the substrate from the atomizer nozzle

- The inclination of the spray cone due to the substrate

- The gas-impinging angle

- The protrusion length of the nozzle

- The shape of the substrate itself

- The movement of the substrate

When considering the above process parameters, the process variables that are

used in assessment of spray rolling process in this study are tabulated in Table 3.2

Table 3.2 Process variables and corresponding abbreviations

Process Variables Corresponding Abbreviation

The diameter of the melt orifice, melt flow Md The cross sectional area of melt tip of tundish Mt The gas pressure of the atomizer nozzle Pgas The distance of the rollers from the atomizer nozzle d Inclination of the spray cone due to the roller i The protrusion length of the nozzle Pr Rolling speed of the twin-roller Sr Thickness of roll gap Gr Super heating of the alloy melt Ts Gas impinging angle α

.

20

3.2.1 Early Spray Rolling Set-ups’ Design Stages

In order to achieve the continuous production of the flat product several spray

rolling set-ups are prepared and experiments are conducted. In the first three

newly designed set-up continuous production of the alloy strip can not be achieved

but valuable practical information related to the spray rolling process and effects

of process parameters is gained.

Figure 3.1 First experimental spray rolling set-up

21

In the first experimental set-up, atomization of the molten alloy is done by a

horizontal gas jet focused on the stream of molten alloy, which is allowed to flow

vertically from the bottom of the induction furnace. The metal is melted in a

separate induction furnace and then poured into the set-up furnace at 800°C. The

design parameters “i, Pr and d” are adjustable by the movable gas jet nozzle. The

atomizer to roller distance “d” is 50 cm. The protrusion length “Pr” is 20cm. In

this set-up continuous atomization of the molten metal cannot be achieved. After

the first experiment it is decided to decrease the roller to atomizer distance “d” and

the need of a new atomizing nozzle is observed. The photograph and simple sketch

of the set-up is shown in Figure 3.1 and Figure 3.4.a respectively.

Figure 3.2 Second experimental spray rolling set-up.

22

In the second experimental set-up, a new atomizer nozzle is used. The nozzle has a

protrusion length “Pr” of 0 cm, gas impinging angle “α” of 15° , melt orifice

diameter of 12mm. The sketch of atomizer nozzle is shown in Figure 3.3. The

atomizer nozzle to roller distance “d” is 10cm. The molten metal is carried from

the furnace to the nozzle with the help of a heat resistant tube made of alumina

fiber. The photograph and sketch of the setup are shown in Figure 3.2 and Figure

3.4.b. respectively.

The continuous atomization cannot be achieved; main problem observed in this

set-up is the early solidification of the molten alloy inside of the heat resistant tube

therefore a heat insulation box is entegrated to the set-up to decrease the heat loss

from the tube. Heat insulation box is made from firebricks fixed inside a metal

box.

Figure 3.3 Sketch of Atomizer Nozzle

23

Figure 3.4 Sketches of experimental set-ups

24

In the third experimental set-up, the heat insulation box that is integrated to the

system, retarded the solidification of the molten metal in the tube. Atomization of

the molten alloy is taken place but it is observed to be inhomogeneous therefore

continuous atomization of the molten alloy is not achieved. The atomizer used in

this set-up produces a high angle spray cone, which result in wide spreading of the

semi solid powders, which is not appropriate for spray deposition through the

roller where a lower angle cone is needed for directing onto the rollers. The heat

insulation box is also decided to be insufficient which cannot avoid the heat loss

from the system successfully. The sketch and photograph of the set-up with the

insulation box is shown in Figure 3.4.c. and Figure 3.5 respectively.

Figure 3.5 Third experimental spray rolling set-up.

25

3.2.2 Advanced Design of Spray Rolling Set-up

When the results of previous three set-ups are assessed; it is decided that one of the

critical process parameter is the superheat of the alloy melt. Although the alloy is

poured at 800°C, the temperature of the alloy melt near to the nozzle is

significantly lower than that as a result the superheat of the alloy is completely

lost. Other critical process parameters are related to the nozzle. The new nozzle

design should accomplish the controlled directing of the alloy melt in the semi

solid powder form and they should hit the rollers just before solidification. It is

concluded that an advanced set-up design concerning the above conclusions is

needed.

First design rationale is the loss of the superheat of the alloy melt which can be

solved by providing external heat to the system and by supporting the molten alloy

front by a heat mass. In order to achieve these goals and electrical resistance

furnace is designed and produced. The electrical furnace includes a tundish unit

having a reservoir capable of keeping 5 liters of molten alloy.

The electrical furnace is designed to reach a temperature of 1000°C and to be able

to keep the molten Al-Fe-V-Si alloy at 800°C - 850°C.

Production of the Electrical Resistance Furnace

The first concern in the furnace design was the insulation of the furnace. For the

insulation 50mm. thick alumina fiber block insulators are used. The block

insulators are designed to withstand 1500°C. A closed box is formed from the

block insulators and then block insulators are carved for the reinforcement of the

electrical resistance wire.

26

For resistance heating, 4 meters of coiled resistance wire is needed to be reinforced

inside of the 4 insulator blocks surrounding the furnace. 20 meters of Nickel-

Chromium 80-20 resistance wire having a diameter of 1.6mm is used, which has a

maximum operating temperature of 1100°C. The resistance wire is coiled into a

outer diameter of 9 mm. The coiled wire is stretched with a ratio of 3:1 as a result

4 meters of coiled electrical resistance wire is obtained.

The electrical resistivity of the 80-20 Nickel-Chromium wire is 108µohm.cm. The

linear resistance of the wire is calculated below;

R (ohm/m) = Resistivity (µohm-cm) / A (in mm²) x 100

= 108µohm.cm / 2,0106mm2 x 100

R = 0,537 ohm/m

For 20m. wire total resistance is 10,74 ohm.

From the experimental observations it is known that the variac unit that is

connected to the furnace can provide 18amperes of current to the 4 m. coiled wire.

Thus the power of the electrical furnace can be calculated as;

P = I2 x R

= (18 amperes)2 x 10,74 ohms

= 3479,76 W 3,4796 kW

The photographs of the production of the electrical furnace is shown in Figure 3.6

27

Figure 3.6 Photographs of electrical resistance furnace production.

28

As a result of electrical resistance furnace production the problem of insufficient

superheat of alloy melt is solved and continuous flow of the allow melt is

achieved. The next design parameter mainly concerned is the nozzle design. In the

new set-up a flat horizontal gas nozzle is used for atomization, which is 25mm

wide and has a inner thickness of 1mm. In the new set-up most of the critical

process parameters can be controlled, using the advanced set-up numerous

experiments are conducted with changing the process variables. Last three set-ups

are classified as Stage 1, Stage 2 and Stage 3 having a positive evolution in the

achievement of the continuous flat product.

Figure 3.7 Advanced experimental spray rolling set-up

29

In design stage 1, electrical resistance furnace is heated up to 800°C. The details of

the process variables are given Table 3.3. After the molten metal that is melted in

the induction furnace is poured into tundish atomization of the melt is begun, as a

result of the large tundish tip the melt flow rate was to high and the atomization of

the melt resulted a flow of semi solid particles rather than semi solid powders.

Powder sized atomization is achieved near the end of batch, while the spray

deposition is taking place it is observed that a significantly large spray cone is

formed and as a result spray deposition has taken place in a wide area over the

rolles so solidification between the roll gap cannot be achieved. After the

experiment it is decided to decrease the roller to atomizer distance to obtain a

smaller spray cone, to decrease the melt flow rate to obtain finer atomization.

In the design stage 2, electrical resistance furnace is heated up to 850°C to obtain a

better super heating of the alloy since the metal flow rate is decreased. After the

atomization is started, it is observed that the solidification between the roll gap is

succeeded and spray rolled products are obtained. However, the production is not

continuous short strips are formed. Subsequent to formation in the roller one set of

strips allowed to cool in water and the others allowed cooling in air. It is decided

that the roller speed should be decreased and the melt flow rate should be

decreased to have better spray deposition over the rollers and better solidification

in the roll gap. Details of process variables is shown in Table 3.3

In the design stage 3, the electrical resistance furnace is heated up to 850°C. A new

composition of Al-Fe-V-Si alloy having lower Si content is used. As a result of

decreased roller speed and homogeneous spray deposition achieved by the

decreased molten metal flow rate, formation of continuous flat product of Al-Fe-

V-Si alloy is achieved. Details of process variables is shown in Table 3.3 The

photograph of flat Al-Fe-V-Si alloy is shown in Figure 3.8

30

Figure 3.8 Sketch showing the process variables

Table 3.3 Detailed tabulation of process variables in the related design

stages

Process Variables & Abbreviation

Design Stage 1 Design Stage 2 Design Stage 3

The cross sectional area of melt tip of tundish

Mt 30mm x 15mm 20mm x10mm 10 mm x 5mm

The gas pressure of the atomizer nozzle

Pgas 98.06 kPa 196.13 kPa 196.13kPa

The distance of the rollers from the atomizer

nozzle d 400mm 300mm 300mm

The protrusion length of the nozzle

Pr 20mm 10mm 10mm

Rolling speed of the twin-roller

Sr 92.3 mm/s 103.3 mm/s 67.0 mm/s

Thickness of roll gap Gr 2mm

3mm end product 2mm

3mm end product 2mm

3mm end product Super heating of the

alloy melt Ts

Furnace kept at 800°C

Furnace kept at 850°C

Furnace kept at 850°C

Gas impinging angle α 0° 0° 0°

31

Figure 3. 9 The continuous flat Al-Fe-V-Si alloy product

Figure 3. 10 Photograph of spray rolling process

32

3.3 Hot Rolling

Hot rolling of the specimens is made in the same twin roller that is used in the

spray-rolling set-up. The specimens are heated up to 450°C in the electrical

resistance furnace and kept for 120 min. before rolling. Hot rolling resulted a %20

reduction in the thickness of the specimen.

3.4 X-Ray Diffraction Study

X-ray diffraction study is made for characterization of phases present in the spray

rolled products. The study is performed by Rigaku International Company

D/MAX2200/PC ULTIMA X-ray Diffractometer by CuKα radiation having

λ:1.5409.

3.5 Optical and Scanning Electron Microscopy Study

Optical and Scanning Electron Microscopy is carried out to examine the

microstructural features of the spray rolled products. The equipment used for SEM

is a JSM-6400 Electron Microscope (JEOL) equipped with NORAN Series II

Microanalyzer System. This microscope is capable of having secondary and

backscattered electron images and EDX microanalysis.

Etching of the Al-Fe-V-Si alloy specimens is done by the Kellers reagent.

33

3.6 Hardness Test

Brinell hardness (HB) values are obtained by using Heckert hardness testing

machine. Eight data is taken from each specimen aiming the assessment of

hardness distribution and calculation of the average hardness.

Hardness covers several properties as resistance to deformation, resistance to

friction and abrasion. The following correlation links brinell hardness with tensile

strength, while resistance to deformation is dependent on modulus of elasticity.

Tensile Strength (MPa) = 3.55 x HB (for HB<175); 3.38 x HB (for HB>175)

3.7 Three Point Bending Test

Three point bending test is done with Shimatsu 10kN testing machine from which

the Force(N) vs Stroke(mm) data is obtained and then Flexural Stress vs Strain

graphs are drawn by the use of following formulations. The maximum flexural

stress values calculated by 0.2% offset method since the bending theory is

applicable only for elastic deformation.

Bending test specimens have span lenght of 35mm, width of 10mm and thickness

of 3.0mm.

Stress σ = M . c / I where c = t/2, M = P . L /4 and I = w . t3 /12 Strain ε = 12 . y .c / L2 where c = t/2 P: Load applied by the testing machine, t: Thickness of the specimen w: Width of the specimen, and L: Span length respectively y: Deflection, stroke applied by the machine.

34

3.8 Tensile Test

Tensile test is done with Shimatsu 10kN testing machine from which the Force(N)

vs Stroke(mm) data is obtained and then Stress (MPa) vs Strain graphs are drawn

and tensile strength and % elongation values are determined.

Tensile test specimens have 8.5mm gauge width and 20mm gauge length with

thickness values of 3.0mm and 2.4mm(hot rolled).

3.9 True Density Measurement Study

True density measurements of the specimens are made by the

Ultrapycnometer1000 Quantachrome Corp., Helium Pycnometer device in the

Central Laboratory of METU.

Conclusions concerning porosity are derived by comparing the measured true

density of the specimens from different processing conditions.

35

CHAPTER 4

RESULTS AND DISCUSSIONS

4.1 Phase Characterization

It is known that the process parameters in production significantly affect the

microstructural and mechanical features of the final product. Since the primary

concern of this study is to design a spray rolling process, the microstructural and

mechanical property results obtained in each process development stage is used for

further design concerns.

Phase characterization study is made to investigate the morphology, existence and

distribution of both expected and unexpected intermetallic and silicide phases in

the spray rolled Al-Fe-V-Si alloy. In order to access the results clearly

metallographic and X-ray specimens are classified due to the processing type and

stage, which is shown in Table 4.1

36

Table 4.1 X-ray and metallographic specimen classification.

Specimen Code Alloy Comp. * Set-Up Design Stage * Further Processing

A1 Al-8Fe-2V-8Si 1 Rolled

A2 Al-8Fe-2V-8Si 2 Water Cooled

A3 Al-8Fe-2V-8Si 2 Air Cooled

B1 Al-8Fe-1V-2Si 3 Water Cooled

B2 Al-8Fe-1V-2Si 3 Air Cooled

B3 Al-8Fe-1V-2Si 3 Hot Rolled

* 1.5%wt Mg is added to all alloys. Design stages are explained in detail at Chapter 2.

4.1.1 Optical and Scanning Electron Microscopy Results

The microstructure of A1 specimen, which is spray deposited and then rolled in as

cast condition features; high level of porosity and starlike precipitates, formerly

characterized as Al3Fe and Al13Fe4 by Sahoo et al. [16,49] , which are finer than

the conventionally cast intermetallics, can be shown in Figure 4.1. Moreover, band

shaped rodlike intermetallic phases are observed in a separate area, which is

evident to the inhomogeneous cooling in the preform. These intermetallic phases

are characterized as Al13Fe4 by X-ray diffraction, and in agreement with the

characterized phases in the previous study by Kalkanli et al. [50] shown in

Figure4.2. It is stated that in spray deposition of Al-Fe-V-Si alloy growth of these

intermetallics are impeded as a result of high cooling rate [50] therefore one can

conclude that the cooling rate achieved in the present setup is not adequate. An

intermetallic phase that is observed is characterized as V3Si by X-ray diffraction, is

unevenly distributed over the Al13Fe4 intermetallics, probably formed as a result of

high Si content in the alloy. EDX analysis is shown in Figure 4.3 and SEM and

optical micrograph is shown in Figure 4.4.

37

(a)

(b)

Figure 4.1 Starlike intermetallics (a) in specimen A1 (b) from Sahoo et al. [10]

38

Figure 4.2 Inhomogeneous area featuring Al13Fe4 intermetallics.

Element

Weight

Conc %

Atom

Conc %

Mg 0.03 0.05

Al 1.52 2.38

Si 23.36 35.23

V 73.74 61.31

Fe 1.35 1.03

Figure 4.3 EDX analysis of V3Si intermetallic.

39

(a)

(b)

Figure 4.4 V3Si intermetallic phase marked by * (a) SEM micrograph in X1000

magnification, (b) optical micrograph in X500 magnification.

40

The specimen A2 is formed in the second design stage of the set-up, which

features smaller tundish opening to decrease molten metal flow aiming to achieve

finer atomization of the stream. In addition to this a water tank is integrated to the

system in order to achieve a higher cooling rate in the as formed metal strip.

The microstructure of the A2 specimen is finer than A1 due to the better

atomization achieved by decreasing the molten metal stream. Large amount of

porosity is observed in the specimen and the phases present are characterized as

band shaped rod like Al13Fe4 intermetallics distributed densely through the matrix

and V3Si intermetallics distributed unevenly through the matrix. No starlike

intermetallics are observed probably the growth is avoided due to the cooling rate

achieved during the atomization. The growth of Al13Fe4 intermetallics into band

shaped morphology is a result of inadequate cooling rate. Although water-cooling

is expected to introduce a higher cooling rate, the amount of water is not enough to

cool the as cast preform it rather be heated up and resulted in an inadequate

cooling rate. Optical and SEM micrographs and EDX analysis of A2 is shown in

Figure 4.5 and Figure 4.6.

(a)

41

(b)

Figure 4.5 (a) optical micrograph of A2 in X50 magnification, (b) SEM

micrograph showing Al13Fe4 and V3Si (*) intermetallics in X300

magnification.

Element

Weight

Conc %

Atom

Conc %

Mg 0.04 0.07

Al 0.88 1.38

Si 23.98 36.17

V 74.31 61.79

Fe 0.79 0.60

Figure 4.6 EDX analysis of V3Si intermetallic in specimen A2.

42

Specimen A3 is produced by the same tundish and that is used with A2, only

difference is A3 is cooled in air after spray rolled into a strip. Microstructure of A3

features homogeneous distribution of rounded intermetallics having an average

diameter of 10µm. There is also band shaped Al13Fe4 intermetallics but they are

observed as cracked. Moreover, V3Si intermetallic phase is observed in the

specimen, which expected due to the high Si content in the alloy composition.

Optical micrographs of specimen A3 showing fine rounded Al13Fe4 intermetallics

is given in Figure 4.7, and SEM micrograph showing round intermetallics, V3Si

intermetallic and cracked band shaped intermetallic is given in Figure 4.8.

Figure 4.7 Optical Micrograph of A3 showing fine rounded Al13Fe4

intermetallics in X500 magnification.

43

(a)

(b)

Figure 4.8 SEM Micrograph of A3 showing (a) fine rounded Al-Fe

intermetallics and V3Si (*) in X400 magnification, (b) cracked Al-Fe

intermetallic in X2500 magnification.

44

Specimen B1 and B2 is produced in the third design stage set-up, which features a

smaller spraying distance and smaller tundish tip opening so that the air

atomization of the alloy and simultaneously rolling of the semi-solid deposit is

achieved, resulting a continuous flat Al-Fe-V-Si alloy product. In addition to this,

Si content in the alloy is decreased hence changing the alloy composition to Al-

8Fe-1.4V-1.8Si + 1.3Mg in %wt.

Specimen B1 is allowed to cool in water after spray rolling and heating up of the

water similar to the case in A2 is observed. In the microstructure of B1 band

shaped Al13Fe4 intermetallics are observed, they are homogeneously distributed in

the matrix and finer when compared to the A2, One can conclude that the finer

microstructure obtained is a result of successful atomization of the melt and

growth of intermetallics to band shape is resulted from the slower cooling rate

during the solidification of the strip. The poor water-cooling is suspected to be

related to the low amount of water used. Since it is seen that adequate cooling rates

can be achieved in air-cooling, instead of using a higher amount of water air-

cooling is chosen for cooling of the spray rolled preform. Optical and SEM

micrographs of B1 are shown in Figure 4.9, 4.10.

45

(a)

(b)

Figure 4.9 (a) Optical micrograph of B1 showing fine intermetallics in X50 magnification. (b) SEM micrograph of B1 showing Al13Fe4 intermetallic phase.

46

Figure 4.10 SEM micrograph of B1 showing the fine dendrite structure

Specimen B2 is produced in the same set-up as B1 so that successful atomization

and the continuous production of the flat product is achieved, moreover B2 is

cooled in air after spray rolling. Microstructure of B2 features fine grains and fine

rounded Al-Fe intermetallics having an average diameter of 3 - 5µm. In addition

an Aluminum rich zone is observed in B2 specimen where there is slight amount

of intermetallics. Kalkanli et al reported a similar zone in previous study, it is told

that Aluminum rich zone is formed during spray deposition as a result of transient

liquid formation and segregation of Fe, V and Si atoms during solute rejection

from that zone [50]. The intermetallic rich regions around the Al rich zone is in

agreement with previous study and probably formed by segregation of solute

atoms. Optical and SEM micrographs of B2 are shown in Figure 4.11, 4.12.

47

(a)

(b)

Figure 4.11 (a) SEM micrograph of B2 showing fine rounded Al-Fe intermetallics in X1000 mag. (b) showing segregation zone in X1000 magnification.

48

(a)

(b)

Figure 4.12 (a) Optical micrograph of B2 showing the Al rich segregation

zone in X100 magnification. (b) EDX analysis taken from the Al rich zone

Element

Weight

Conc %

Atom

Conc %

Mg 0.60 0.69

Al 88.87 92.91

Si 1.99 2.00

V 1.54 0.85

Fe 7.00 3.54

49

Specimen B3 is produced in the same set-up as B2. For the further processing of

B3 specimen a hot rolling stage is introduced, aiming the consolidation of the flat

product, which contains considerable amount of porosity. Prior to hot rolling

specimen is kept at 450°C for 120mins, since the alloy contains Magnesium

precipitation of Mg2Si is considered. In the SEM micrographs of B3 very fine

white particles are observed, EDX analysis taken from these particles show the

presence of Mg together with the other alloying elements (Fe, V, Si). These

particle are presumed to be Mg2Si precipitates or Al13 (Fe, V)3Si silicides, for the

clear characterization of these particles a transmission electron microscopy study

is needed since the particles are much smaller than 1 µm.

Microstructure of B3 features both rounded and band shaped intermetallics and

they are observed as very fine and homogeneously distributed through the matrix,

band shaped Al13Fe4 intermetallics are cracked, porosity level is decreased

considerable only microporosity is observed. SEM and optical micrographs of B3

are shown in Figure 4.12 and 4.13.

50

(a)

(b)

Figure 4.12 Optical micrographs of B3 showing the homogeneous

distribution of intermetallics (a) in X50 (b) X500 magnification.

51

(a)

(b)

Figure 4.13 SEM micrographs of B3; (a) showing the cracked Al13Fe4

intermetallics, (b) showing possible Mg2Si or Al13 (Fe, V)3Si fine white

dispersoids.

52

An interesting microstructure is observed in the cross sections of the flat products

of both B2 and B3. Both microstructures show high intermetallic phase density

through the centerline of the cross section probably due to effect of segregation

during solidification between the rolls. SEM and optical micrographs of this region

revealed an interesting phase different that Al-Fe intermetallics. EDX analysis of

the phase showed a high Mg and V content, the intermetallic phase is

characterized as V2Mg3Al18 by X-ray Diffraction methods. Optical and SEM

micrographs and EDX analysis is given in Figure 4.14, 4.15.

Element

Weight

Conc %

Atom

Conc %

Mg 5.37 6.31

Al 80.82 85.51

Si 0.96 0.98

V 12.85 7.20

Figure 4.14 EDX analysis of V2Mg3Al18 intermetallic phase

53

(a)

(b)

Figure 4.15 (a) Backscattered SEM image of B3 specimen showing

V2Mg3Al18 intermetallic phase. (b) Typical optical micrograph of cross

section of B2 and B3 specimens showing V2Mg3Al18* and Al13Fe4+

intermetallics in X500 magnification.

54

4.1.2 X-Ray Diffraction Results

X-ray diffraction study is done for characterization of the suspected phases in

specimens obtained from each set-up design stage since the changes in processing

parameters affect the existence of secondary phases.

Figure 4.16 X-ray diffractogram of the specimen produced in the 1st design stage,

showing characterized phases; Al13(Fe, V)3Si, Al13Fe4 and V3Si.

X-ray Diff. Of Stage 1

0

2000

4000

6000

8000

10000

12000

14000

20 30 40 50 60 70 80 90 100

2theta

Inte

nsit

y

Phases

Al13(Fe,V)3Si

Al13Fe4

V3Si

55

Figure 4.17 X-ray diffractogram of the specimen produced in the 2nd design stage,

showing characterized phases; Al13(Fe, V)3Si, Al13Fe4 and V3Si.

In the specimens of the first and second design stage the phases that are

characterized are; the bcc Al13(Fe, V)3Si silicide phase having the lattice

parameter a: 1.260 nm, the monoclinic Al13Fe4 phase having the lattice parameters

as a: 1.548nm b: 8.083nm c: 1.247nm and the cubic V3Si phase having the lattice

parameter of a: 0.472nm.

X-Ray Diff. Of Stage 2

0

500

1000

1500

2000

2500

3000

3500

4000

4500

5000

20 30 40 50 60 70 80 90 100

2theta

Inte

nsit

y

Phases

Al13(Fe,V)3Si

Al13Fe4

V3Si

56

Figure 4.18 X-ray diffractogram of the specimen produced in the 3rd design stage,

showing characterized phases; Al13(Fe, V)3Si, Al13Fe4, V2Mg3Al18 and V3Si.

X-ray Diff. Of Stage 3

0

2000

4000

6000

8000

10000

12000

14000

16000

18000

20 30 40 50 60 70 80 90 100

2theta

Inte

nsit

y

Phases

Al13(Fe,V)3Si

Al13Fe4

V3Si

V2Mg3Al18

57

Figure 4.19 X-ray diffractogram of the specimen produced in the 3rd design stage

and hot-rolled, showing characterized phases; Al13(Fe, V)3Si, Al13Fe4, V2Mg3Al18

V3Si and Mg2Si.

One can conclude from the above X-ray diffractograms that the amount of V3Si is

decreased in the 3rd design stage and it does not exist in the hot-rolled specimen.

Moreover a new phase is characterized in the diffractograms of 3rd stage and hot

rolled specimen as V2Mg3Al18 , which is in agreement with the SEM study results.

X-ray diffractogram of the hot rolled specimen also shows the possible presence of

the Mg2Si phase.

X-ray Diff. Of Stage 3 Hot-Rolled

0

2000

4000

6000

8000

10000

12000

14000

16000

18000

20000

20 30 40 50 60 70 80 90 100

2theta

Inte

nsit

y

Phases

Al13(Fe,V)3Si

Al13Fe4

V3Si

V2Mg3Al18

Mg2Si

58

4.2 Mechanical Test Results

4.2.1 Three Point Bending Test Results

Three point bending test is applied to the specimens from the design stage that the

continuous flat products are obtained, which are designated as B1 and B2 before.

Maximum bending stress is measured by 0.002 offset method since the bending

theory is applicable only in elastic deformation. Flexural strength values obtained

are tabulated below in Table 4.2.

Table 4.2 Three point bending test results of B2 and B3.

Specimen Result1 (MPa) Result2 (MPa) Result3 (MPa) Maximum Flexural

Strength (MPa) B1 (watercooled) 161.24 207.84 227.11 198.73 B2 (air cooled) 266.26 271.83 254.54 264.21

Theoretical Tensile Strength of the Al-Fe-V-Si alloy is 462 MPa [1].

Flexural strength values are below the expected theoretical level, this probably

because of the considerable amount of porosity and the inhomogenities present in

the specimens, discussed in microscopic examination section.

59

One can conclude from the results in Table 4.2 that the flexural strength of B2 is

significantly higher than that of B1. When the microstructures of these specimens

are concerned B1 features fine and rod like band shaped Al13Fe4 intermetallics

while in B2 these intermetallics have become rounded as a result of the cooling

rate achieved. Therefore, it can be said that growth of Al13Fe4 intermetallics into

rod like band shaped morphology degrades mechanical properties so critical care

should be given to the control of intermetallic morphology by the verification of

cooling rate.

60

4.2.2 Tensile Test Results

Tensile test is applied to specimens from the continuously formed air cooled B2

and the continuously formed air cooled and further hot rolled B3. All tested

specimens are cracked in the elastic region and no plastic deformation is observed

when theoretical tensile behavior of the alloy is concerned it is seen that Al-Fe-V-

Si alloy should represent a plastic deformation and ultimate tensile strength value

about 462 MPa. [1]. Therefore, the results obtained in this study can be examined

as yield strength values while these values are still significantly below the

theoretical values 413 MPa.[1]. This insignificance in tensile strength test values is

probably caused by poor specimen preperation. Although significant tensile

strength values are not reached, elongation results are still valuable for

comparison. As it can be seen from the below figure %elongation value thus

ductility is decreased after hot rolling. The lower strength value that is observed in

hot-rolled specimen is unexpected and probably caused because of the cracking of

the intermetallics. The tensile test results are given in Table 4.3. One of the stress-

strain diagrams is given in the Appendix B.

Table 4.3 Tensile strength and %elongation values of B2 and B3

Specimen Result 1 (MPa -

%elongation)

Result 2 (MPa -

%elongation)

Result 3 (MPa -

%elongation)

Average Fracture Strength (MPa)

Average %

elongation

B2 (Air-cooled)

162.72 – 6.35 107.54 – 5.95 122.48 – 6.66 130.91 6.32

B3 (Hot- rolled)

96.32 – 4.09 83.27 – 3.79 91.35 – 4.06 90.31 3.98

Theoretical Tensile Strength of the Al-Fe-V-Si alloy is 462 MPa [1].

61

4.2.3 Hardness Test Results

Hardness test of the specimens from all previously classified design stages such

that A1, A2, A3, B1, B2, B3, is conducted by using the Brinell Hardness Scale

(HB). Ultimate tensile strength (UTS) is also estimated from the average Brinell

hardness values. Hardness test results are given in Table 4.4.

Table 4.4 Brinell Hardness test results.

Specimen A1 A2 A3 B1 B2 B3

1 125 112 92 102 102 92

2 89 116 89 121 71 92

3 102 125 125 109 105 89

4 130 130 89 116 99 92

5 121 140 112 140 102 92

6 102 135 99 145 102 92

7 121 125 125 99 99 92

8 116 121 116 109 105 89

Average Hardness

(HB)

113.25 125.50 105.88 117.63 98.13 91.25

Standart Deviation

(HB)

13.17 8.76 14.50 15.83 10.47 1.30

Average UTS

(MPa)

402.04 445.53 375.86 417.57 348.34 323.94

Standart Deviation

(MPa)

46.75 31.10 65.48 56.20 47.17 4.61

Theoretical Tensile Strength of the Al-Fe-V-Si alloy is 462 MPa [1].

62

Ultimate tensile strength values obtained from the hardness test are closer to the

theoretical values contrary to the tensile and three point bending results. Therefore,

one can conclude that by the spray rolling set-up designed in this study, aimed

mechanical strength values are achieved by microstructural development but

improvements in macro-scale is still needed to be made in the further studies.

Examination of the hardness results revealed that the A2 and B1 specimens, which

are cooled in water, have higher hardness values than A3 and B2 specimens that

are cooled in air. When the microstructures are considered they feature greater

band shaped intermetallics instead of round intermetallics in A3 and B2, which

contributed to hardness. However, it should be considered that the band shaped

intermetallics caused early fracture in bending test.

When hardness distribution is taken into account, the uneven distribution is

observed in A1, A2, A3 and B1 specimens. In B2 specimen the homogeneity of

hardness and microstructure is observed with an exception of the transient liquid

phase region, discussed in microscopical results section also featured a low

hardness value of 71 HB. It is clear that this region can be avoided by control of

process parameters, since in hot rolled B3 specimen hardness and microstructure is

observed to become more homogeneous which can be derived from the decrease in

the standard deviation in the hardness. The decrease in the hardness of the hot

rolled specimen is probably due to the cracking of the intermetallics that are

observed in the SEM studies.

63

4.3 True Density Measurement Results

True densities of two specimens from B1 (spray rolled) and B3 (hot rolled after

spray rolling) are measured by helium pycnometer and compared in Table 4.5.

Table 4.5 True Density Measurement Results

Specimen Temperature (°C) Weight (g) Volume (cc) Density (g/cc)

B1-a 25.9 1.1824 0.3787 3.1223

B1-b 25.9 1.2685 0.4129 3.0724

B3-a 26.1 1.2423 0.3752 3.3115

B3-b 26.1 1.3224 0.3848 3.4362

According to the measured true density results, it can be seen that the density is

increased after hot rolling, therefore one can conclude that the porosity is

decreased after hot rolling of the spray rolled product.

64

CHAPTER 5

CONCLUSIONS

1. It is observed that molten metal flow rate directly affects the atomization of

the melt, concurrently the refinement of the microstructure is dependent on

the atomization, and therefore the refinement of the microstructure depends

on the atomization conditions before the spray rolled product is solidified

between the rolls.

2. The cooling rate just after the deposition over the rollers and during the

solidification of the spray rolled strip, significantly affects the growth rate

of the Al13Fe4 intermetallics hence their morphologies.

3. In low cooling rates monoclinic Al13Fe4 intermetallics having rod like band

shaped morphologies are formed. The larger intermetallics contribute to

hardness values while it is observed that they degrade the flexural strength,

when the cooling rate is increased they become finer and rounded showing

an increase in the flexural strength.

4. Microstructural features are directly affected by the change in process

parameters, which contribute to the processing conditions. Therefore

microstructural evaluation is a useful tool for assessing process parameters.

65

5. SEM EDX and XRD study of the spray rolled Al-8Fe-1.7V-8Si alloy

revealed an extraordinary phase characterized as V3Si having rounded

morphology, which is formed as a result of the high Si content in the alloy.

6. SEM EDX and XRD study of the spray rolled Al-8Fe-1.4V-1.8Si alloy

revealed a extraordinary phase characterized as V2Mg3Al18 having rounded

morphology and an average size of 10µm.

7. Hot rolling of the spray rolled specimens resulted in a homogeneous

microstructure thus a homogeneous hardness distribution and a significant

decrease in the porosity.

8. In the further hot rolling of the spray rolled flat product, band shaped

Al13Fe4 intermetallics are cracked, resulting a decrease in hardness values

and degrading the mechanical properties. It is concluded that hot rolling

parameters should be revised.

9. One of the spray rolled products exhibit an Aluminum rich zone featuring

very slight amount of intermetallics. It is formed during spray rolling as a

result of transient liquid formation and segregation of Fe, V and Si atoms

during solute rejection from that zone. The hardness of the zone is

measured to be 71HB, which is very low compared to the rest of the

product.

66

CHAPTER 6

SUGGESTIONS FOR FUTURE WORK

1. The as-spray rolled flat product formed in the developed spray rolling set-up

exhibit the desired microstructural features and homogeneous distribution of

the second phase particles. However, a considerable amount of porosity is

observed and it is also shown that it could be removed by applying further hot

rolling. Another drawback is the bended structure of the as-spray rolled alloy

sheet. Integration of a hot twin rolling stage to the present set-up, in which the

alloy strip is allowed to be rolled just after production is believed to enhance

the properties of the end product.

2. The integrated further hot rolling stage will allow the deformation of the

product just after solidification, therefore it may avoid the cracking of the

intermetallics since the deformation is applied during the growth of the

intermetallic phases.

3. In this study, it is observed that when different cooling mediums i.e. water, air;

are introduced just after the production of the spray rolled strip, different

microstructures are obtained. Therefore, it is suspected that the cooling rate

after spray rolling may affect the growth of the intermetallics. For the further

studies, it is suggested that this conclusion could be examined by measuring

the temperature of the as-spray rolled strip and comparing it with the

decomposition temperature data of the alloy which could be obtained from

DSC study.

67

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72

APPENDIX A

A : X-RAY DIFFRACTION CARDS OF THE CHARACTERIZED PHASES

Table A.1 X-ray card of Aluminum

73

Table A.2 X-ray card of Al13(Fe,V)3Si

74

Table A.3 X-ray card of Al13Fe4

75

Table A.4 X-ray card of Mg2Si

76

Table A.5 X-ray card of V3Si

77

Table A.6 X-ray card of V2Mg3Al18

78

APPENDIX B

B: SAMPLE STRESS VS STRAIN GRAPH OF B2 SPECIMEN

Table B.1 Stress ve Strain graph of B2 specimen

Stress vs Strain Graph of B2

162,72

0,00

20,00

40,00

60,00

80,00

100,00

120,00

140,00

160,00

180,00

0 0,01 0,02 0,03 0,04 0,05 0,06 0,07

Strain

Str

es

s (

MP

a)